Coherent Nanodispersion-Strengthened Shape-Memory Alloys

ABSTRACT

High-strength, low-hysteresis TiNi-based shape-memory alloys (SMAs) employing fully coherent low-misfit nanoscale precipitates, wherein the precipitate phase is based on an optimized composition for high parent-phase strength and martensite phase stability, and compensating the stored elastic energy through the addition of martensite stabilizers. The alloys, with a yield strength in excess of 1200 MPa, are useful for applications such as self-expanding stents, automotive actuators, and other applications wherein SMAs with high output force and long cyclic life are desired.

CROSS REFERENCE TO RELATED APPLICATIONS

This is a continuation utility application based upon U.S. Ser. No.10/809,082, filed Mar. 25, 2004, which is based upon provisionalapplication U.S. Ser. No. 60/457,418, filed Mar. 25, 2003, entitled“High Performance Shape Memory Alloy,” which is incorporated herein byreference in its entirety and for which priority is claimed.

BACKGROUND OF THE INVENTION

In a principal aspect, the present invention relates to high-strength,low-hysteresis shape-memory alloys (SMAs), and in particular TiNi-basedSMAs, employing coherent, low-misfit nanoscale-sized precipitates. Suchalloys are contemplated to have a myriad of practical applicationsincluding, but not limited to, use in medical stents and actuators.

The shape-memory effect is a consequence of a crystallographicreversible, thermoelastic martensitic transformation. SMAs rely onproperty changes induced during the transformation from a hightemperature phase (parent phase) to a low temperature phase (productphase or martensite); the product phase is relatively compliant incomparison with the parent phase. Shape-memory actuation occurs when anSMA is deformed in its martensite state, below its M_(s) temperature;the deformed shape is maintained upon unloading. Once reheated beyondthe austenite finish temperature (A_(f)), an SMA will work against aresisting force to regain its original shape.

Superelasticity occurs when an SMA is deformed above austenite starttemperature (A_(s)), but below M_(s) ^(σ) (the highest temperaturepossible to have martensite). In this range, martensite can be madestable with the application of stress, but becomes unstable again whenthe stress is removed. Because of superelasticity, SMAs can deformelastically up to large strains and recover perfectly without beingdamaged by unloading, similar to rubber.

Under constrained conditions, the output stress of an SMA duringreversion of martensitic transformation is typically limited by the flowstrength of the parent phase. For engineering applications, it is alsohighly desirable, if not essential, that the shape-memory behavior isrepeatable and predictable after many cycles through the transformation.Therefore, to improve both the output force and the cyclic lifetime ofSMAs, the strength of the alloy (i.e. flow strength of parent phase)must be improved. By raising the critical shear stress for slip, theirreversible slip deformation during the martensite reorientation andstress-induced martensite transformation can be suppressed, which, inturn, improves the shape-memory effect and transformationsuperelasticity characteristics.

Currently, the three most commonly used SMAs in engineering applicationsare TiNi and the copper-based alloys, CuZnAl and CuAlNi. Severaliron-based SMAs, such as FeMnSi, FePt, and FePd are also the subjectmatter of research for industrial applications. However, TiNi-basedalloys are currently the most widely used SMAs due to good corrosionresistance and biocompatibility.

Various types of precipitate strengthening may be considered in theTiNi-based system. On the Ti-rich side of the binary Ti—Ni system, aTi₂Ni dispersion can be obtained while on the Ni-rich side Ni₃Ti/Ni₄Ti₃precipitates can be considered for strengthening dispersions. Kajiwaraet al. [Philos. Mag. Lett., 1996, vol. 74, pp. 137-144, J. Phys. IV,2001, vol. 11, pp. 395-405, and Metall. Mater. Trans. A, 1997, vol. 28,pp. 1985-1991 (incorporated herewith)] found that subnanometric thinplate metastable bct precipitates formed when sputter-deposited Ti-richTiNi shape-memory films are annealed in the temperature range of 377 to827° C. Due to the relatively low heat treatment temperature, diffusionof Ti atoms is not rapid enough to form stable Ti₂Ni precipitates;instead, Guinier-Preston zone-type precipitates which contain excess Tiatoms are produced. With these fine precipitates in the parent phase,they could achieve recovery strength 670 MPa. However, theseprecipitates have been observed only after annealing of sputterdeposited and amorphous TiNi thin films.

While precipitation strengthening has been mainly considered inTiNi-base thin films or bulk single crystals, deformation processing hasbeen considered in bulk polycrystalline TiNi alloys for strengthening.Lin and Wu [Acta Metall. Mater., 1994, vol. 42, pp. 1623-1630(incorporated herewith)] studied cold-rolled equiatomic TiNi alloys.With a cold rolling at room temperature to the extent of 31% reductionin thickness, they improved the yield stress from 380 MPa of a solutiontreated specimen to 1000 MPa. However this is an economicallyinefficient approach, as the alloys have to be heat-treated followingeach cold rolling step.

Koizumi et al. [Mater. Sci. Engng A.: 1997, vol. 223, pp. 36-41(incorporated herewith)] examined the high-temperature strength of TiNialloys in the context of developing new alloys to replace Ni-basesuperalloys. They demonstrated that a dispersion of Heusler phase(Ni₂TiAl-type with L2₁ structure) increases the compressive yieldstrength of 50.7Ni—40.9Ti—8.4Al (in at %) by an order of magnitude up to2300 MPa. This strengthening method is potentially applicable to boththin film and bulk alloy processing. While they have achieved impressivecompressive yield strength in the TiNi-based alloys with Heuslerprecipitates, they did not consider any of the shape-memorycharacteristics of the alloys. Their alloys were solely developed ashigh-temperature materials, neglecting the thermoelastic transformationsand the superelasticity.

In CuZnAl-based SMAs, a ductile second phase cc can be distributedwithin the β matrix. Compared with that of the single-phase alloy, thefatigue life of dual-phase alloys with homogeneously distributedglobular α phase is increased in both the martensitic and superelasticstate [Metall. Trans. A, 1992, vol. 23, pp. 2939-2941 (incorporatedherewith)]. Semi-coherent γ precipitates in the β matrix have beenstudied by Lovey and Cesari [Mater. Sci. Engng A.: 1990, vol. 129, pp.127-133 (incorporated herewith)]. In CuAlNi-based SMAs, theprecipitation of coherent γ₂ intermetallic compound (Cu₉Al₄) can beconsidered [J. Phys. IV, 1995, vol. 5 (C2), pp. 193-197 (incorporatedherewith)]. In FeMnSi-based SMAs, the addition of small amounts of Nband C is known to produce very small NbC carbide precipitates inaustenite, which improves the shape memory effect [Scripta Mater., 2001,vol. 44, pp. 2809-2814 (incorporated herewith)]. Various phases usefulfor strengthening the parent phase of the matrix are summarized inTABLE 1. TABLE 1 SMA Parent Phase Strengthening Phases B2-TiNiMetastable Ti₂Ni, Ni₄Ti₃, Stable Ni₃Ti, L1₂ (Ni₂TiAl) β-CuZnAl α, γβ-CuAlNi γ₂ (Cu₉Al₄) FeMnSi NbC

Nonetheless, microstructural design and the implementation in processingfor improving the strength of SMAs while controlling the transformationtemperatures have remained a scientific and engineering challenge.High-strength, low-hysteresis TiNi-based SMAs as well as other SMAswhich achieve yield strength greater than 1200 MPa while maintainingdesired transformation temperatures are much needed.

SUMMARY OF THE INVENTION

Briefly, the invention comprises high-strength, low-hysteresis SMAs and,in particular TiNi-based SMAs, employing coherent low-misfit nanoscalesize precipitates, wherein the precipitate phase is based on anoptimized composition for high parent-phase strength and martensitephase stability, utilizing martensite stabilizers to compensate for thestored elastic strain energy. Cycled TiNi alloys frequently exhibitdecreased recovery forces and recoverable strain, all the while showingincreased permanent strain and shifts in the transformationtemperatures. To improve the output force and the cyclic lifetime ofTiNi-based alloys, the strength of the parent phase can be significantlyimproved by appropriate additions of nanodispersions through alloyingelements such as Al and an additive selected from the group consistingof Zr, Hf, Pd, Pt and combinations thereof.

More broadly, SMAs may achieve increased strength and thereby highoutput force as well as long cyclic life without irreversible effects bythe addition of additives, which provide for low misfit between therespective phases (i.e. additives which result in coherency). Suchadditives preferably produce less than about 2.5% misfit in the latticeparameter.

Another feature of this invention comprises the ability to predictivelycontrol the phase transformation temperatures and to minimize hysteresisas a result of the low misfit. The additives are martensite stabilizers,which compensate for the elastic energy stored in the non-transforming,coherent, nanodispersion. While Zr is a preferred additive (along withAl) in the TiNi system, other additives such as Hf, Pd and Pt orcombinations thereof are useful.

Additionally, the technique of matching phases within the parametersdisclosed may be applied to other SMAs including but not limited toCuZnAl, CuZnNi, iron-based SMAs and various TiNi-based SMAs. Accordingto known lattice constants [Pearson's Handbook of Crystallographic Datafor Intermetallic Phases, ASTM International, Newbury, Ohio, 1991(incorporated herewith)], the misfit between α and β in CuZnAl-basedSMAs is about 21%, misfit between γ and β is about 0.7%, and inCuAlNi-based SMAs, the misfit between γ₂ and β is about 0.9%. InFeMnSi-based SMAs, the crystal structure of NbC compound is of the NaCltype and its lattice constant is 0.4470 nm, larger by 24% than thelattice constant of the austenite (fcc) 0.3604 nm. Additives preferablyproduce a misfit less than the values listed above, while providingtransformation temperature control.

Thus, it is an object of the invention to provide a new class of SMAsthat can achieve a yield strength greater than 1200 MPa whilemaintaining desired transformation temperatures.

Another object of the invention is to provide high-strength,low-hysteresis SMAs employing coherent low-misfit nanoscale sizeprecipitates wherein the interphase misfit is less than about 2.5%.

Yet another object of the invention is to provide high-strength,low-hysteresis SMAs with long-term microstructural cyclic stability,wherein the fatigue life is greater than about 10 million cycles.

Another object of the invention is to provide TiNi-basednanodispersion-strengthened alloys wherein the microstructure comprisescoherent low-misfit nanoscale size precipitates.

A further object of the invention is to provide TiNi-basednanodispersion-strengthened alloys wherein the microstructure comprisescoherent low-misfit multicomponent Heusler nanodispersions distributedin the parent phase.

Another object of the invention is to provide composition tolerance byincorporating a third multicomponent phase as a buffer for excess Ti inthe nanodispersion-strengthened TiNi-based SMA.

A further object of the invention is to provide composition tolerance byincorporating a bcc β Nb—Ti phase as a buffer for excess Ti in thenanodispersion-strengthened TiNi-based SMA.

These and other objects, advantages and features will be set forth inthe detailed description which follows.

BRIEF DESCRIPTION OF THE DRAWINGS

In the detailed description that follows, reference will be made to thedrawings comprising the following figures:

FIG. 1 is a flow block logic diagram that characterizes the designconcepts of the alloys of the invention;

FIG. 2 is an equilibrium phase diagram depicting the phases andcomposition at various temperatures in the pseudo-binary TiNi—NiAlsystem relative to the preferred embodiment and example of theinvention;

FIG. 3 is a graph showing the solution temperature vs. Zr content in apreferred embodiment and example of the invention;

FIG. 4 is a graph showing the partitioning of Zr between B2-TiNi andL2₁-Heusler phases in a preferred embodiment and example of theinvention;

FIG. 5 is a graph showing ambient interphase lattice misfit vs. Zrcontent in a preferred embodiment;

FIG. 6 is a schematic showing cross-sectional drawings of aTiNi-actuated microvalve in the a) closed and b) open positions, as anexemplary application of the invention;

FIG. 7A is a TEM dark-field micrograph showing coherent nanoscalecuboidal Heusler precipitates in an example of an alloy of theinvention, Ni—45Ti—5Al (in at %) aged at 600° C. for 2000 h;

FIG. 7B is a TEM dark-field micrograph showing coherent nanoscalespheroidal Heusler precipitates in an example of an alloy of theinvention, Ni—40Ti—5Al—5Zr (in at %) specimen aged at 600° C. for 2000h; and

FIG. 8 is a graph showing the compressive stress-strain response at roomtemperature of an embodiment of the invention, Ni—47Ti—3Al—25Pd (in at%) aged at 600° C. for 100 h.

DETAILED DESCRIPTION OF THE INVENTION

FIG. 1 is a systems flow-block diagram which illustrates theprocessing/structure/properties/performance relationships for alloys ofthe invention. The desired performance for the application (e.g.self-expanding stent, microactuators in microelectromechanical systems,SMA patch repair, etc.) determines a set of alloy properties required.Alloys of the invention exhibit the structural characteristics that canachieve the desired combination of properties and can be assessedthrough the sequential processing steps shown on the left of FIG. 1.

Employing the concepts reflected by FIG. 1, following are the criteriafor the physical properties and the microstructure and compositioncharacteristics for the alloys. This is followed by the processabilitycharacteristics of the alloys, applications, the experimental resultsrelating to the discovery and examples of the alloys that define, ingeneral, the range and extent of the elements, physical characteristicsand processing features of the present invention.

Physical Characteristics

The physical characteristics or properties of the most preferredembodiments of the invention are generally as follows:

Strength equivalent to or better than cold-worked SMA, i.e.:

-   -   Yield Strength≧1200 MPa.    -   Fatigue life longer than 10 million cycles.    -   Optimum microstructural features for transformation temperatures        and maximum output strength/fatigue resistance.        Microstructure and Composition Characteristics

The alloy designs achieve improved output force and cyclic lifetime viananoscale, coherent, low-misfit precipitates without causingirreversible effects on the martensitic transformation. Lattice misfitarising from different lattice parameters between two coherent phasescauses coherency strains with an associated volume strain energy thatcan act as obstacles to martensite interfacial motion, potentiallyincreasing the transformation hysteresis (A_(f)-M_(s)). The hysteresisof the martensitic transformation determines the response rate of thefinal application. A quantitative theory for such behavior has beendeveloped by Grujicic, Olson and Owen [Metall. Trans. A, 1985, vol. 16,pp. 1713-1722] which is incorporated herewith. In a system withplastically deforming precipitates, the hysteresis width will increaseif these particles do not participate in the transformation.Irreversible plastic deformation of a particle will contribute to theinterfacial friction stress as the interface intersects it. In NiTiNballoys, the irreversible deformation of the Nb-rich phase delays therecovery, increasing the hysteresis. 47Ni—44Ti—9Nb (in at %) is acommercially used alloy exhibiting a wide transformation temperaturehysteresis, useful for coupling and sealing. Widening of the hysteresishas also been observed in CuZnAl SMAs, where plastic accommodationoccurred in γ type precipitates due to matrix shape change upontransformation.

In contrast, as discovered in the subject invention, SMAs strengthenedby coherent, low-misfit, nanoscale precipitates show no significantincrease in transformation hysteresis, indicating no significantinterfacial friction from the precipitates. The coherent, low-misfitprecipitates lower the chemical equilibrium To temperature, which is thetemperature at which the parent and martensite have the same Gibbs freeenergy. For precipitate particles of equilibrium phases which do nottransform into martensite, but are elastically sheared by thetransformation, a significant amount of reversible elastic strain energyis stored. This stored energy is equivalent to further undercooling. Thechemical driving force due to the undercooling is given by ΔG=Δs ΔTwhere Δs is the entropy change of the transformation per unit volume,and ΔT=T₀−T is the amount of undercooling from the chemical equilibriumtemperature T₀.

Another potential effect of the coherency strains is the loss ofcoherency of precipitates by cycling through the transformation. Thishas been observed in a nonthermoelastic FeNiC martensite by Chen andWinchell [Metall. Trans. A, 1980, vol. 11, pp. 1333-1339] which isincorporated herewith. Since shape-memory-based devices are typicallycycled many times, to ensure long-term microstructural cyclic stability,the interphase misfit has to be reduced. Thus, a feature of the alloysof the invention is minimization of the interphase lattice mismatch topromote fine scale homogeneous precipitation.

The strength of an overaged material is inversely proportional to anaverage particle spacing, or it scales with √{square root over (ƒ)}/rwhere ƒ is the phase fraction and r is the particle size. Therefore fora given phase fraction, the finest and closely spaced dispersion ofstrengthening particles is desired. This can be achieved by increasingthe thermodynamic driving force for nucleation, which, in turn, isachieved by increasing the supersaturation or reducing the latticemisfit.

The precipitation of equilibrium Heusler (Ni₂TiAl-type with L2₁structure) phase in TiNi is useful to satisfy the design criteria andtherefore is considered as a preferred embodiment of the subjectinvention. There is a lattice misfit between TiNi and Ni₂TiAl, asdetermined by the relation$\delta = {( \frac{a_{{Ni}_{2}{TiAl}} - {2a_{TiNi}}}{2a_{TiNi}} ) = {- 0.0257}}$where α_(Ni) ₂ _(TiAl) is the ambient lattice parameter of Ni₂TiAl(a=0.5865 nm) and α_(TiNi) is the ambient lattice parameter of TiNi(a=0.3010 nm). The misfit decreases at elevated temperatures due to acombined effect of solute solubility limit and thermal expansions, andtherefore about 2.5% is the upper limit for a tolerable misfit in thesubject invention. The lowest possible misfit between B2 and L2₁ phasescan be achieved by increasing the lattice parameter of themulticomponent Heusler phase through alloying elements such as Hf or Zrsubstituting on the Ti sublattice, and Pd or Pt substituting on the Nisublattice in the alloy.

Al added to form the Heusler phase has significant solubility in the B2matrix. Al dissolved in the matrix also decreases the transformationtemperatures drastically. While transformation temperatures arerelatively insensitive to the Ni/Ti ratio in the Ti-rich regime, theyshow a strong decrease in the Ni-rich regime. Since Al is substitutingin the Ti sublattices, the transformation temperature is affected bothby the overall atomic percentage of Al as well as the adjusted Ni/Tiratio. Because of the strong decrease of transformation temperatures byAl in B2, elements which can stabilize the martensite phase and therebyoffset the B2 stabilizing effect of soluble Al are added. AccordinglyHf, Zr, Pd, and Pt, initially considered for reducing the lattice misfitbetween B2 and L2₁ phase, are also martensite stabilizers. Theiraddition allows a higher transformation temperature. If Hf, Zr, Pd, andPt partition to B2, the stability of martensite phase will be increased,and if they partition of L2₁, the interphase lattice misfit will bereduced.

For comparison, in both CuZnAl and CuAlNi-based systems, the stabilityof β′ martensite decreases with Al, Zn, and Ni content. The martensitictransformation temperature is very sensitive to small variations inalloy composition. Although the transformation temperatures of bothCuZnAl and CuAlNi alloys can be manipulated over a wide range, thepractical upper limits are 120° C. and 200° C. respectively, above thesetemperatures the transformations tend to be unstable. FeMnSi basedalloys are one-way shape memory materials with high strength, highaction temperatures, good workability and low cost. Addition of nitrogenor rare earth elements lowers the M_(s) temperature, stabilizing theaustenite after shape recovery.

In TiNi-based SMAs, oxides such as Ti₄Ni₂O or Y₂O₃ can form during thearc-melting or powder consolidation process; however such dispersionsmay be desirable because of their grain refining effect. Typical TiNicontains oxygen concentrations of 350 to 500 ppm and carbon from 100 to500 ppm depending on starting materials and melt practice. Ti₄Ni₂O typeoxides effectively pin the grain boundaries during the dynamicrecrystallization occurring with the hot-working process. To improve theductility of the material the grain size has to be reduced, and for thispurpose B is preferably added to form borides. Yang and Mikkola [ScriptaMetall. Mater., 1993, vol. 28, pp. 161-165 (incorporated herewith)],confirmed improved ductility by the addition of 0.12 at % boron inTiNiPd alloys.

Another feature of the alloys is built-in tolerance for compositionvariation to ensure a robust design. The composition range of theB2-TiNi phase is narrow even at high temperatures. Therefore, strictcomposition control in alloy production would be required to avoidprecipitation of Ni₃Ti, Ni₄Ti₃, Ni₃Ti₂, or Ti₂Ni that are harmful toductility. To promote robust alloy production, the composition tolerancein manufacturing will have to be increased. Ni-rich compositions areavoided because the martensitic transformation temperatures dramaticallydecrease. By incorporating a third multicomponent phase as a buffer forexcess Ti in the nanodispersion-strengthened TiNi-based SMA, tolerancefor composition variance can be built in. For example, the bcc β Nb—Tiphase can be incorporated as a buffer for excess Ti in alloycompositions that are deliberately kept Ni-lean to avoid the competingNi-rich phases. Variations in excess Ti would be absorbed in smallcomposition variations in the Nb-based buffer phase, which is kept at asufficiently low phase fraction not to degrade mechanical properties andtransformation hysteresis. To prevent increasing of the hysteresis, asseen in the commercially used NiTiNb alloys, Nb will have to be keptlower than about 9 at %.

Processability Characteristics

A principal goal of the subject invention is to provide alloys with theobjective physical properties and microstructural characteristicsrecited above and with processability that renders the alloys useful andpractical. With a number of possible processing paths associated withthe scale of manufacture and the resulting cleanliness and quality for agiven application, compatibility of the alloys of the subject inventionwith a wide range of processes is desirable and is thus a feature of theinvention.

A primary objective and characteristic of the alloys is compatibilitywith melting practices such as Vacuum Induction Melting (VIM) and VacuumArc Remelting (VAR), and other variants such as Vacuum Induction SkullMelting process. Alloys of the subject invention can also be produced byother processes such as powder consolidation. By selection ofappropriate elemental content in the alloys of the subject invention,the variation of composition can be minimized.

Allowable variation results in an alloy that can be homogenized atcommercially feasible temperatures, usually at metal temperatures inexcess of 900° C. Objectives regarding solution heat treatment includethe goal to fully homogenize the alloy while maintaining a fine scalegrain refining dispersion (i.e. Ti₄Ni₂O, Y₂O₃) and a small grain size.The solution temperature of binary TiNi shape memory alloy is generallylimited by the order-disorder transition temperature at 1090° C. With aninitial target solution temperature of 900° C., the Al content of thematrix can be designed utilizing a pseudo-binary phase diagram, FIG. 2,of TiNi to NiAl. This is created using the thermodynamic calculationsoftware Thermo-Calc [Calphad, 1978, vol. 2, pp. 227-238 (incorporatedherewith)] and a custom thermodynamic database, which is based on athermodynamic assessment in the Ti—Ni—Al system undertaken incollaboration with Dr. Weiming Huang at QuesTek Innovations, LLC [Jung,J. Doctoral Thesis, Department of Materials Science and Engineering,Northwestern University, Evanston, Ill., 2003 (incorporated herewith)].This could also be constructed empirically by a person skilled in theart. From this, the solubility of aluminum at 900° C. can be determinedas about 6 at %.

In a preferred embodiment, a Zr addition to the TiNi-Heusler system isconsidered because Zr has the most significant effect on decreasing thelattice misfit while efficiently raising the martensite stability.Adding small amounts of Zr increases the solution temperature as seen inFIG. 3, which is again calculated using the custom database (referencedabove) with Thermo-Calc. This could also be assessed empirically by aperson skilled in the art. Since Zr quickly increases the solutiontemperature, it was determined the Al level should preferably be lowerthan about 4 at %.

Due to the solution hardening of solute atoms in the TiNi B2 matrix, thealloys are stronger than binary TiNi even before the strengthening phaseprecipitation. This can make manufacturing and machining difficult,since for these operations a soft material which exhibits favorable chipformation is desired. Therefore alloys of the subject invention arepreferably annealed to reduce the hardness before they are supplied to amanufacturer. Typically this pretreatment would be accomplished byheating the alloy at about 800° C., for a period of less than onethousand hours, preferably between one and one hundred hours and coolingto room temperature. In some cases a multiple-step annealing process mayprovide more optimal result. In such a process an alloy of the inventionmay be annealed at a series of temperatures for various times that mayor may not be separated by an intermediate cooling step or steps.Through this pretreatment, alloys would be over-aged to coarsenprecipitates and reduce the alloying elements in the B2 matrix, therebyminimizing solid solution strengthening. Components made of alloys ofthe subject invention can be manufactured or machined after thispretreatment, and the components will be ultimately given a finalsolutionizing and aging treatment to attain full hardening.

The temperature of the final aging process would typically be between600° C. and 800° C., at a temperature where the lowest possible misfitcan be achieved by increasing the lattice parameter of themulticomponent Heusler phase. To realize this design concept, acombination of Analytical Electron Microscopy and 3-Dimensional AtomProbe microanalysis was conducted by Jung et al. [Metall. Mater. Trans.,2003, vol. 34, pp. 1221-1235 (incorporated herewith)] and the interphasepartitioning at 600° C. and 800° C. were established. The B2/L2₁ solutepartitioning is discovered to be strongly temperature dependent and canreverse direction between 600° C. and 800° C. By incorporating thesemeasurements into a solution thermodynamic assessment for theTi—Ni—Al—Zr system, the composition dependence of the solutepartitioning can be predicted, and a model for the composition andtemperature dependence of the B2 and L2₁ lattice parameters has beendeveloped to predictively control interphase misfit at precipitation anduse temperatures.

For the temperature dependent partitioning of Zr, FIG. 4 can begenerated. The partition coefficient of Zr shows a smooth compositiondependence, in addition to the temperature dependence. In thesecalculations the Al content of the alloy was kept at about 5 at %. Thesolute partitioning of Zr is discovered to be in favor of reducing theinterphase misfit at 600° C. Therefore the modeling efforts are focusedat 600° C., and better agreement is obtained at this temperature. MoreZr enters the Heusler phase at low Zr contents at 600° C. Using modelsfor the composition and temperature dependence of the B2-TiNi andL2₁-Heusler phase lattice parameters, the misfit of the B2 and Heuslerphases at 600° C. can be plotted as a function of Zr content, in thealloy, as seen in FIG. 5. Thus the temperature of the final agingprocess would preferably be from 600° C. to 650° C. and less thanhundred hours in duration, preferably between one and twenty hours. Theoutcome of the desired process is a B2 matrix strengthened by afully-coherent low-misfit nanoscale dispersion which is aged at aminimum predetermined temperature for a minimum time to achieveworkability.

Applications

Among many, a few applications could be considered to test the limits ofthe conceptual design capability for coherentnanodispersion-strengthened shape-memory alloys.

Medical applications such as self-expanding stents utilize thesuperelasticity of TiNi-based SMAs, for which the T₀ will have to beplaced below body temperature. The biased stiffness of TiNi causes thestent to passively press against the vessel in a very compliant fashion,yet the stent resists constriction with a comparatively high stiffness.Physicians can oversize the stent to the vessel, and feel confident thatwhile the stent is stiff enough to scaffold the vessel, the passiveforces will not be so great as to perforate the vessel wall. To improvethe cyclic lifetime of TiNi, the strength of the alloy parent phase mustbe improved to eliminate accommodation slip during transformation, whichcan be achieved through the subject invention.

Recently, the demand for powerful microactuators inmicroelectromechanical systems (MEMS) has motivated significant SMA thinfilm research. This is because SMAs produce actuation forces and strokessuperior to other actuator materials. As seen in FIG. 6, in theoff/unpowered position the microspring deflects the martensitic TiNifilm downward, pressing the boss against the orifice opening. Whenheated, the austenitic TiNi becomes nearly flat, deflecting themicrospring upward, lifting the boss away from the orifice and allowingfluid to flow. Traditional SMA microactuators used in MEMS devicessuffer from limited cyclic life due to accommodation slip. To improvethe output force and the cyclic lifetime of TiNi-based alloys, thestrength of the alloy must be improved.

Shape-memory actuators are becoming increasingly popular for automotiveapplications. In a modern car more than 100 actuators are used tocontrol engine, transmission and suspension performance, to improvesafety and reliability and enhance driver comfort. The operatingtemperature range of a car ranges from −40° C. to approximately +100°C., with even higher temperatures in under-hood locations. In order towork properly at all temperatures, the shape memory alloy has to have anM_(f) temperature well above the maximum operating temperature.

A novel technique can be developed based on the subject invention thatapplies a self-repair patch across cracked weld joints so thatcatastrophic fatigue failures could be prevented. Welded structures thatare subjected to cyclic loading often fail by fatigue at the weld joint.This can lead to the structure eventually breaking or becomingnon-functional. In either case the cost associated with the fatiguefailure can be significant. Repair of cracked or aging structures withbonded composite patches has shown great promise to become a viablemethod for life extension of such structures. This process relies on theprinciple of crack-closure phenomenon where the opening stress on thecrack faces is reduced by placing a patch across the wake of the crack.However these patch designs merely act as a band-aid to hinder futurecrack growth. The shape memory mechanism of TiNi-based alloys can beutilized to bear loads and apply compressive stress to the crack. Apre-deformed shape memory alloy patch can be heated above the austenitefinish temperature and the patch will apply crack closure clamping forceby reverting to its memorized shape.

Tailored to the applications as described above, the T₀ temperature canbe calculated and alloy compositions can be designed accordingly takinginto account the effect of stored elastic energy of precipitates.

EXPERIMENTAL RESULTS AND EXAMPLES

A series of prototype alloys were prepared. Sample buttons or slugs ofprototype alloys weighing 25 g were prepared by arc-melting in an argonatmosphere using pure elements (99.99˜99.994 wt % Ni, 99.99 wt % Ti,99.999 wt % Al, 99.9 wt % Hf, 99.98 wt % Pd, 99.95 wt % Pt, and 99.999wt % Zr). Taking equiatomic TiNi as reference, in alloys A, A+5Hf andA+5Zr the Ni-content was kept at 50 at % while Ti was partially replacedby Al, Hf or Zr. On the other hand, in alloys B+5Pd, B+20Pd and B+5Pt,Ni was partially substituted by Pd and Pt. Alloys with high Pd-contentNi—49Ti—1Al—25Pd (D+1Al) and Ni—47Ti—3Al—25Pd (D+3Al) were alsodesigned. A prototype alloy of composition Ni—32Ti—3Al—15Zr (E+15Zr) wasinvestigated. The compositions of the prototypes are given in TABLE 2.TABLE 2 Alloy Ni Ti Al Hf Zr Pd Pt A 50 45 5 — — — — A + 5Hf 50 40 5 5 —— — A + 5Zr 50 40 5 —  5 — — B + 5Pd 45 44 6 — —  5 — B + 20Pd 30 44 6 —— 20 — B + 5Pt 45 44 6 — — — 5 D + 1Al 25 49 1 — — 25 — D + 3Al 25 47 3— — 25 — E + 15Zr 50 32 3 — 15 — —Note:All values in at %

Consistent with model predictions, A+5Zr prototype alloy demonstratesnear-zero misfit at 600° C. As this alloy was designed to stabilize B2against martensitic transformation, the martensitic transformationtemperature was too low (<−150° C.) to be detected. The Al retained inthe B2 matrix decreases the transformation temperature, while themartensite stabilizer Zr was present only in a limited amount.

To study the multicomponent effect on T₀, transformable alloys with highPd-content D+1Al, D+3Al, and high Zr-content E+15Zr were designed. Allthree second-iteration prototypes gave satisfactory transformationtemperatures. The best mechanical properties of the second iterationwere exhibited by alloy E+15Zr, which demonstrates a recovery stress of2100 MPa at 180° C., in combination with a high A_(f) reversiontemperature of 149° C.

Following is a summary of the described experiments and alloys:

Example 1

As-cast specimen of alloy A in TABLE 2 was sealed in an evacuated quartzcapsule and solution treated at 1100° C. for 100 h. After quenching bycrushing the capsules in oil, it was annealed at 800° C. for 1000 h or600° C. for 1000 or 2000 h in evacuated quartz capsules, and thenquenched into oil. FIG. 7A shows coherent nanoscale cuboidal Heuslerprecipitates in alloy A aged at 600° C. for 2000 h.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 3together with Vickers hardness numbers. Originally intended for thephase-relations study, this alloy was over-aged to yield large Heuslerprecipitates. Therefore the effect of precipitation strengthening isminimized and the hardness numbers mainly reflect the solutionstrengthening contribution.

As this alloy was designed to stabilize B2 against martensitictransformation, the martensitic transformation temperatures were too low(<−150° C.) to be detected. Hot hardness measurements were carried outon the aged A alloys using a custom tester. Both A alloys aged at 800°C. and 600° C. exhibit a monotonic decrease of hardness over thetemperature. This indicates that the formation of stress-inducedmartensite is suppressed and the room temperature hardness measurementsreflect the strength of the parent phase, i.e. M_(s) ^(σ)<roomtemperature. TABLE 3 Various measured properties for Alloy A SolutionProperty Treated Aged at 800° C. Aged at 600° C. Lattice 0.30018 nm0.30022 nm (B2) 0.30132 nm (B2) Parameter (B2) 0.59068 nm (Heusler)0.59358 nm (Heusler) Vickers 349 463 430 Hardness Number

Example 2

As-cast specimen of alloy A+Hf in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 1100° C. for 100 h. Afterquenching by crushing the capsules in oil, it was annealed at 800° C.for 1000 h or 600° C. for 1000 or 2000 h in evacuated quartz capsules,and then quenched into oil.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 4.Substitution of Ti by Hf leads to an increase in lattice parameter ofthe quaternary alloys, compared to alloy A. Vickers hardness numbers forA+Hf aged at 800° C. or 600° C. for 1000 h are also shown in TABLE 4.Originally intended for the phase-relations study, this alloy wasover-aged to yield large Heusler precipitates. Therefore the effect ofprecipitation strengthening is minimized and the hardness numbers mainlyreflect the solution strengthening contribution. As this alloy wasdesigned to stabilize B2 against martensitic transformation, themartensitic transformation temperatures were too low (<−150° C.) to bedetected. TABLE 4 Various measured properties for Alloy A + 5Hf SolutionProperty Treated Aged at 800° C. Aged at 600° C. Lattice 0.30331 nm0.30298 nm (B2) 0.30171 nm (B2) Parameter (B2) 0.59410 nm (Heusler)0.59518 nm (Heusler) Vickers — 491 489 Hardness Number

Example 3

As-cast specimen of alloy A+Zr in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 1100° C. for 100 h. Afterquenching by crushing the capsules in oil, it was annealed at 800° C.for 1000 h or 600° C. for 1000 or 2000 h in evacuated quartz capsules,and then quenched into oil.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 5.Substitution of Ti by Zr leads to an increase in lattice parameter ofthe quaternary alloys, compared to alloy A. A+5Zr demonstrates near-zeromisfit at 600° C. This is consistent with the Heusler particlemorphology transition to a spheroidal form, as shown in FIG. 7B.

Vickers hardness numbers for A+Zr aged at 800° C. or 600° C. for 1000 hare shown in TABLE 5. Originally intended for the phase-relations study,this alloy was over-aged to yield large Heusler precipitates. Thereforethe effect of precipitation strengthening is minimized and the hardnessnumbers mainly reflect the solution strengthening contribution. As thisalloy was designed to stabilize B2 against martensitic transformation,the martensitic transformation temperatures were too low (<−150° C.) tobe detected. TABLE 5 Various measured properties for Alloy A + 5ZrSolution Property Treated Aged at 800° C. Aged at 600° C. Lattice0.30543 nm 0.30468 nm (B2) 0.30255 nm (B2) Parameter (B2) 0.59851 nm(Heusler) 0.60351 nm (Heusler) Vickers 503 522 491 Hardness Number

As the A+5Zr prototype showed promising interphase misfit levels, theprecipitation strengthening was investigated in detail. A+5Zr was agedat 600° C. for 1, 3, 10, and 100 h and Vickers hardness was measured asa function of aging time. The measured properties are listed in TABLE 6.The average equivalent spherical radius of the precipitates wasdetermined based on conventional transmission electron microscopymeasurements. Peak hardening is in the range from 1 to 10 h of aging at600° C., which corresponds to a precipitate radius of 1.44 to 2.45 nm.TABLE 6 Vickers hardness number of A + 5Zr aged at 600° C. Aging TimePrecipitate Radius Vickers Hardness Solution — 503 1 hour 1.44 nm 516 3hours 1.86 nm 553 10 hours 2.45 nm 531 100 hours 4.37 nm 531 1000 hours5.65 nm 491 2000 hours 9.94 nm 404

Example 4

As-cast specimen of alloy B+5Pd in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 1100° C. for 100 h. Afterquenching by crushing the capsules in oil, it was annealed at 800° C. or600° C. for 100 h in evacuated quartz capsules, and then quenched intooil.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 7.Substitution of Ni by Pd leads to an increase in lattice parameter ofthe quaternary alloys, compared to alloy A. As this alloy was designedto stabilize B2 against martensitic transformation, the martensitictransformation temperatures were too low (<−150° C.) to be detected.TABLE 7 Lattice Parameter for Alloy B + 5Pd Heat Treatment LatticeParameter Aged at 800° C. 0.30360 nm (B2) 0.60175 nm (Heusler) Aged at600° C. 0.30276 nm (B2) 0.59818 nm (Heusler)

Example 5

As-cast specimen of alloy B+20Pd in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 1100° C. for 100 h. Afterquenching by crushing the capsules in oil, it was annealed at 800° C. or600° C. for 100 h in evacuated quartz capsules, and then quenched intooil.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 8.Substitution of Ni by Pd leads to an increase in lattice parameter ofthe quaternary alloys, compared to alloy A. As this alloy was designedto stabilize B2 against martensitic transformation, the martensitictransformation temperatures were too low (<−150° C.) to be detected.TABLE 8 Lattice Parameter for Alloy B + 20Pd Heat Treatment LatticeParameter Aged at 800° C. 0.30612 nm (B2) 0.61031 nm (Heusler) Aged at600° C. 0.30579 nm (B2) 0.60397 nm (Heusler)

Example 6

As-cast specimen of alloy B+5Pt in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 1100° C. for 100 h. Afterquenching by crushing the capsules in oil, it was annealed at 800° C. or600° C. for 100 h in evacuated quartz capsules, and then quenched intooil.

The ambient lattice parameters obtained from the X-ray diffractionexperiments, corrected for instrumental factors, are listed in TABLE 9.Substitution of Ni by Pt leads to an increase in lattice parameter ofthe quaternary alloys, compared to alloy A. As this alloy was designedto stabilize B2 against martensitic transformation, the martensitictransformation temperatures were too low (<−150° C.) to be detected.TABLE 9 Lattice Parameter for Alloy B + 5Pt Heat Treatment LatticeParameter Solution Treated 0.30059 nm (B2) Aged at 800° C. 0.30612 nm(B2) 0.61031 nm (Heusler) Aged at 600° C. 0.30579 nm (B2) 0.60397 nm(Heusler)

Example 7

As-cast specimen of alloy D+1Al in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 950° C. for 100 h. A lowsolutionizing temperature was chosen to minimize the possible nucleationand growth of Ti₂Ni-based particles. This alloy was designed to studythe multicomponent effect on T₀ in a single phase material, andtherefore was not aged. The A_(s), A_(f), M_(s), and M_(f) temperatureswere determined by Differential Scanning Calorimetry (DSC). These aresummarized in TABLE 10. TABLE 10 Martensitic Transformation Temperaturesfor Alloy D + 1Al M_(s) 135° C. M_(f) 119° C. A_(s) 130° C. A_(f) 145°C. Hysteresis (A_(f) − M_(s))  10° C.

Example 8

As-cast specimen of alloy D+3Al in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 950° C. for 100 h. A lowsolutionizing temperature was chosen to minimize the possible nucleationand growth of Ti₂Ni-based particles. After quenching by crushing thecapsules in oil, it was annealed at 600° C. for 100 h in evacuatedquartz capsules, and then quenched into oil. The A_(s), A_(f), M_(s),and M_(f) temperatures were determined by Differential ScanningCalorimetry (DSC). These are summarized in TABLE 11. The large decreasein transformation temperatures, and increased hysteresis compared toD+1Al indicate the effects of Al. Based on the solubility limit of Al inTiNi at 600° C., D+3Al should yield 2.4% Heusler phase by volume. TABLE11 Martensitic Transformation Temperatures for Alloy D + 3AlSolutionized Aged M_(s) −34° C. −23° C. M_(f) −86° C. −37° C. As −47° C.−21° C. A_(f) −10° C.  −2° C. Hysteresis (A_(f) − M_(s))  24° C.  21° C.

Compression tests specimens of 3 (diameter)×5 (height) mm were producedby EDM machining. The aged D+3Al shows a superelastic behavior, at thestress level up to 700 MPa, as shown in FIG. 8. D+3Al aged at 600° C.exhibits a strength level comparable to precipitation strengthenedbinary TiNi alloys, which is encouraging considering the Heusler phasefraction is only 2.4%.

Example 9

As-cast specimen of alloy E+15Zr in TABLE 2 was sealed in an evacuatedquartz capsule and solution treated at 950° C. for 100 h. A lowsolutionizing temperature was chosen to minimize the possible nucleationand growth of Ti₂Ni-based particles. After quenching by crushing thecapsules in oil, it was annealed at 600° C. for 100 h in evacuatedquartz capsules, and then quenched into oil. The A_(s), A_(f), M_(s),and M_(f) temperatures were determined by Differential ScanningCalorimetry (DSC). These are summarized in TABLE 12. Thedispersion-strengthened E+15Zr shows no significant increase intransformation hysteresis, indicating no significant interfacialfriction from the precipitates. TABLE 12 Martensitic TransformationTemperatures for Alloy E + 15Zr Solutionized Aged M_(s) 137° C. 143° C.M_(f) 104° C. 112° C. A_(s) 127° C. 128° C. A_(f) 142° C. 149° C.Hysteresis (A_(f) − M_(s))  5° C.  6° C.

Compression tests specimens of 3 (diameter)×5 (height) mm were producedby EDM machining. Aged E+15Zr specimens were compressed above A_(f) at180° C. and 155° C. For the first specimen tested at 180° C., yield wasaround 2100 MPa and fracture occurred at 2200 MPa. After observing thefracture, another sample was tested at the same temperature, to checkthe reproducibility of the test. Superelastic behavior was observed at astress level up to 1400 MPa. This is a very high stress level,especially considering a small predicted Heusler phase fraction. Basedon a calculated thermodynamic equilibrium, the volume fraction of L2₁phase is 11.1%. The measured properties are listed in TABLE 13. TABLE 13Compression Test Results for Alloy E + 15Zr aged at 600° C. PropertyValue Yield Strength 1099 MPa at 155° C. 2100 MPa at 180° C. FractureStress 2200 MPa at 180° C.

As a consequence of such research and examples, the alloys in thepreferred embodiment of the subject invention are considered to have arange of combinations of elements as set forth in TABLE 14. TABLE 14Alloy Subclass Ti Al Hf Zr Pd Pt 1 32 to 40 3 to 4 — 8 to 15 — — 2 30 to40 3 to 4 9 to 17 — — — 2 About 47 About 3 — — 5 to 20 — 3 About 47About 3 — — — 5 to 20All values in at %

With one or more of: Nb B O C <9 <0.1 <500 ppm <500 ppmAnd the balance Ni

Preferably, impurities are avoided; however, some impurities andincidental elements are tolerated and within the scope of the invention.Thus, by weight, most preferably, O is less than about 0.05% and C lessthan about 0.05%. Ni-rich compositions should be avoided to prevent theformation of metastable phases such as Ni₃Ti₂ or Ni₄Ti₃. In the Ni-leanregion the low-melting Laves phase should be avoided. To achieve this,the sum of Ti, Al, Hf, and Zr, and the sum of Ni, Pd, and Pt, arepreferably kept at about 50 at %.

The TiNi-based alloys comprise a structure of multicomponent Heuslerphase nanodispersions distributed in the parent phase, wherein theHeusler phase is based on an optimized composition for high parent-phasestrength and martensite phase stability, and compensating the storedelastic energy by the addition of martensite stabilizers. The alloycomposition allows for slight composition variations that may ariseduring processing, by incorporating a bcc Nb—Ti phase as a buffer forexcess Ti in alloy compositions.

This alloy will have to be solutionized at a temperature higher than890° C. and subsequently annealed at about 800° C. between one and onehundred hours to reduce the hardness before they are supplied to amanufacturer. After this pretreatment, the components will be ultimatelygiven a final solutionizing and aging treatment to attain fullhardening. Final aging treatment will be at about 600° C. for 20 h or atabout 650° C. for a shorter time, for peak strength. This design isrobust for aging, because the size evolution of L2₁ precipitates isrelatively slow.

The specific alloy compositions of TABLE 14 represent the presentlyknown preferred and optimal formulations in this class of alloys, itbeing understood that variations of formulations consistent with thephysical properties described, the processing steps and within theranges disclosed as well as equivalents are within the scope of theinvention. Subclass 1 is similar in composition to alloys A, A+5Zr, andE+15Zr of TABLE 2 and is optimal for reducing the lattice misfit whilestabilizing the martensite phase. Subclass 2 is similar in compositionto alloy A+5Hf, and is optimal for reducing the lattice misfit whilestabilizing the martensite phase. Subclasses 3 and 4 are similar incomposition to alloys B+5Pd, B+20Pd, B+5Pt, D+1Al, and D+3Al of TABLE 2and are optimal for superelastic applications.

The subject invention can be extended to other systems of SMAs,including copper-based alloys CuZnAl, CuAlNi, and iron-based SMAs suchas FeMnSi. In CuZnAl-based SMAs, the additive will have to optimize thestrength and phase stability of the β parent phase by reducing themisfit between the parent and strengthening phases such as α or γ, whilecompensating for the stored elastic energy. In CuAlNi-based SMAs, γ₂intermetallic compound (Cu₉Al₄) can be considered as a strengtheningphase. In FeMnSi-based SMAs, strengthening particles such as NbCcarbides will have to be coherently precipitated in a nanoscale-sizewhile maintaining desired transformation temperatures.

While examples of the alloys of the invention, their processing,manufacture and use have been set forth, multiple variations of suchSMAs are considered to be within the scope of the invention. Therefore,the invention including the class of coherentnanodispersion-strengthened SMAs and the processes for making and usingsuch alloys is to be limited only by the following claims andequivalents thereof.

1. A shape memory alloy comprising, in combination: a temperaturesensitive alloy characterized by a displacive transformation between afirst parent phase and a second product phase, said first parent phasemaintaining a deformed shape below the M_(s) temperature followingstress and unloading and transformable to an original shape uponreheating above an A_(f) temperature; said alloy further characterizedby a coherent, nanodispersion of an additional phase providing a misfitof less than about 2.5% in the lattice structure between thenanodispersion and the parent phase.
 2. The alloy of claim 1 comprisinga CuZnAl combination having a β parent phase and a nanodispersionselected from the group consisting of α or δ phase.
 3. The alloy ofclaim 1 comprising a CuAlNi combination and a δ₂ intermetallicnanodispersion.
 4. The alloy of claim 1 comprising a FeMnSi combinationand a NbC nanodispersion.
 5. The alloy of claim 1 comprising a NiTiAlcombination with a Huesler nanodispersion in a B2 parent phase.